Multi-state memory and multi-functional devices comprising magnetoplastic or magnetoelastic materials

ABSTRACT

Apparatus and methods are disclosed that enable writing data on, and reading data of, multi-state elements having greater than two states. The elements may be made of magnetoplastic and/or magnetoelastic materials, including, for example, magnetic shape-memory alloy or other materials that couple magnetic and crystallographic states. The writing process is preferably conducted through the application of a magnetic field and/or a mechanical action. The reading process is preferably conducted through atomic-force microscopy, magnetic-force microscopy, spin-polarized electrons, magneto-optical Kerr effect, optical interferometry or other methods, or other methods/effects. The multifunctionality (crystallographic, magnetic, and shape states each representing a functionality) of the multi-state elements allows for simultaneous operations including read&amp;write, sense&amp;indicate, and sense&amp;control. Embodiments of the invention may be used, for example, for storing, modifying, and accessing data for device, sensor, actuator, logic and memory applications. Embodiments may be particularly effective for non-volatile memory or other read&amp;write, sense&amp;indicate, and/or sense&amp;control functions in computer or other applications; such simultaneous operation of two (or more) of said multiple functionalities open new pathways for miniaturization of devices.

This application has priority of U.S. Provisional Application60/859,163, filed Nov. 14, 2006, which is incorporated herein byreference

FIELD OF THE INVENTION

The invention relates to multi-state memory and/or multi-functionaldevices using magnetoplastic and/or magnetoelastic materials whichcouple magnetic and crystallographic/shape states. More specifically,the present invention relates to crystallographic and magnetic states ofmagnetoplastic and/or magnetoelastic materials, including magnetic shapememory alloys (MSMA) and other materials, to store, access, and modifydata for sensor, actuator, logic and memory applications.

Current conventional logic, memory, and sensing and control technologyis based on binary elements. Typically, a binary element containssemiconductor material that has two states, a conducting one, and aninsulating one. On the other hand, the present invention relates todevices and methods wherein elements made of magnetoplastic and/ormagnetoelastic materials are used, said elements preferably having sixcrystallographic states, and wherein the crystallographic states arelinked to three magnetic states. This enables the combined usage of asingle element as memory and logical unit. As a consequence, thelow-level programming of computers becomes much more efficient. Becausea single element may have six states (rather than two), memory densityincreases 3 fold. Because each of the six states are stable without theuse of an external force (e.g., without external electric field or othercontinued force), the memory provided by said element can benonvolatile.

Magnetoplastic materials, including ferromagnetic shape-memory alloyswith twinned martensite, tend to deform upon the application of amagnetic field (Ullakko 1996, Murray et al. 2000, Chernenko et al.2000). The magnetic-field-induced deformation can be irreversible(magnetoplasticity, Mullner et al. 2002, Mullner et al. 2003a) orreversible (magnetoelasticity, Chernenko et al. 2000, Ullakko et al.1996). The magnetoplastic effect is related to themagnetic-field-induced displacement of twin boundaries, in anirreversible process. The magnetoelastic effect is also related to themagnetic-field-induced displacement of twin boundaries, but in a processthat is at least somewhat reversible. While a strict definition of“elastic” would imply that magnetoelastic materials return withouthysteresis to their initial state after removal of the magnetic-field,the term “magnetoelastic,” as it is frequently used, may includematerials that deform and return to their initial state upon removal ofthe magnetic field either without hysteresis or with some hysteresis.

While literature on the subject of magnetoplasticity andmagnetoelasticity has indicated such materials to be relevant asactuators for converting electrical energy, or changes in magneticfield, to mechanical motion, the invention utilizes, in combination,magnetoplasticity/magnetoelasticity and the reverse or inverse effect(strain-induced change of magnetization) for data reading and writing,which may take the forms of memory and retrieval, displacement andsensing, sensing and control, or sensing and indicating, for example.

The inventors believe that magnetoplastic and/or magnetoelasticmaterials are uniquely suitable as multi-state data storage devices dueto their large range of deformation, small threshold stress andsignificant change of magnetization. See Mullner et al 2003b, Straka andHeczko 2006, Suorsa et al. 2004, below, for discussions of range ofdeformation, threshold stress, and change of magnetization formagnetoplastic materials.

Closed-loop control of any automated process, including machines androbots, contains three main elements: a motor/actuator that drives thetool, a sensor which detects the position of the tool, and a controllerwhich activates the motor. Sensor and controller are linked throughelaborate electronics. In view of miniaturization, the electronicsbecome a limiting aspect of the system. Multifunctional multi-statedevices, according to embodiments of the invention, are a pathway toovercome these limitations by inherently linking them between magneticand crystallographic and/or shape states wherein one function utilizingsaid multi-states will operate as a sensing means and the other functionutilizing said multi-states will operate as control signal means.Similarly, displace&sense, sense&indicate, and read&write operations(such as memory storage and retrieval) are sets of functions that can beperformed sequentially or simultaneously by causing/urging/actuating thematerial to enter various different magnetic and crystallographic and/orshape states (herein called “writing”), and then sensing/responding tosaid various different magnetic and crystallographic and/or shape states(herein called “reading”).

RELATED ART

Background references that may be relevant to embodiments of theinvention include:

Chernenko V A, Cesari E, Kokorin V V, Vitenko I N, Scripta Metal Mater1995; 33:1239.

Chernenko V A, L'vov VA, Cesari E, J Magn Magn Mater 1999; 196-197:859.

Chernenko V A, L'vov VA, Pasquale M, Besseghini S, Sasso C, Polenur D A,Int J Appl Electromag Mech 2000; 12:3.

Chernenko V A, Müllner P, Wollgarten M, Pons J, Kostorz G, J de Phys IV, 2003; 112:951.

Ferreira P J, Vander Sande J B, Scripta Mater 1999; 41:117.

Ge Y, Heczko O, S{umlaut over (R)}derberg O, Lindroos V K, J. Appl. Phys2004; 96: 2159.

Greer J R, Oliver W C, Nix W D, Acta Mater 2005; 53:1821.

Jääskeläinen A, Ullakko K, Lindroos V K, J de Phys IV, 2003; 112:1005.

Murray S J, Marioni M, Allen S M, O'Handley R C, Lograsso T A, Appl PhysLett 2000a; 77:886.

Murray S J, Marioni M, Kulda A M, Robinson J, O'Handley R C, Allen S M,J Appl Phys 2000b; 87:5774.

Müllner P, Int J Mater Res (Z f Metallk) 2006;97:205.

Müllner P, Chernenko V A, Wollgarten M, Kostorz G, J Appl Phys2002;92:6708.

Müllner P, Chernenko V A, Kostorz G, J Magn Magn Mater 2003a;267:325.

Müllner P, Chernenko V A, Kostorz G. Scripta Mater 2003b;49:129.

Müllner P, Chernenko V A, Kostorz G, Mater Sci Eng A 2004a;387:965.

Müllner P, Chernenko V A, Kostorz G, J Appl Phys 2004b;95:1531.

Müllner P, Mukherji D, Erni R, Kostorz G, Proc Int Conf “Solid to solidphase transformations in inorganic materials” PTM'05, Phoenix, Ariz.,Mai 29-Jun. 3, 2005, Vol. 2:171.

Müllner P, Ullakko K, Phys Stat Sol (b) 1998;208:R1.

Pan Q, James R D, J Appl Phys 2000;87:4702.

Pond R C, Celotto S, Intern Mater rev 2003;48:225.

Sozinov A, Likhachev A A, Lanska N, Ullakko K, Appl Phys Lett2002;80:1746.

Sozinov A, Ullakko K, IEEE Trans Magn 2002;38:2814.

Straka L, Heczko O, Scripta Mater 2006;54:1549.

Sullivan M R, Chopra H D, Phys Rev B 2004;70:094427.

Suorsa I, Pagounis E, Ullakko K, Appl. Phys. Lett. 2004a;23:4658.

Suorsa I, Tellinen J, Ullakko K, Pagounis E, J Appl Phys 2004b;95:8054.

Tickle R, James R D, J Magn Magn Mater 1999;195:627.

Ullakko K, J Mater Eng Perf, 1996;5:405.

Ullakko K, Huang J K, Kantner C, O'Handley R C, Kokorin V V, J Appl Phys1996;69:1966.

Ullakko K, Aaltio I, www.adaptamat.com, “Actuator 2002”, Bremen,Germany.

Vlasova N I, Kandaurova G S, Shchegoleva N N, J Magn Magn Mater2000;222:138.

SUMMARY OF THE INVENTION

Apparatus and methods are proposed that comprise writing and reading ofmulti-state “bits” or multi-state elements on magnetoplastic and/ormagnetoelastic materials or other materials that have greater than twomagnetic, crystallographic, and/or shape states. Said “bits” or“elements” refer to individual/single portions/regions of material eachadapted for receiving and holding information (writable) and also forsaid information being read (readable), and wherein “bit” is notintended to refer to traditional binary coding. “Multi-state” refers togreater than two states, and the multi-state bits/elements of thepreferred embodiments are formed or defined by the magnetic andcrystallographic/shape states of magnetoplastic and/or magnetoelasticmaterials Preferred embodiments of the invention may achievemulti-functionality by the multiple crystallographic, magnetic, andshape states of the bits/elements of the magnetoplastic/magnetoelasticmaterials, which allows for simultaneous operations includingread&write, sense&indicate, and sense&control. Simultaneous operation oftwo (or more) functionalities open new pathways for miniaturization ofdevices.

The preferred writing process is conducted through the application of amagnetic field and/or a mechanical action to said magnetoplastic and/ormagnetoelastic materials. It has been shown that switchingmagnetic/crystallographic domains changes the orientation of the surfaceof these materials. The surface inclination can be detected using lightof any wave length and by using other forms of radiation. Thus, lightand other forms of radiation may be used for the reading process. Forexample, the reading process may be conducted through atomic-forcemicroscopy, magnetic-force microscopy, spin-polarized electrons,magneto-optical Kerr effect, optical interferometry or other methods, ora combination of the above.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a schematic depiction of the experimental device with whichdeformation-induced change of magnetization was demonstrated.

FIG. 2 is a graph of stress and magnetization of a Ni₅₁Mn₂₈Ga₂₁ crystalas a function of compressive deformation in the device depicted in FIG.1.

FIGS. 3 (a), (b), and (c) is a schematic depiction of the magnetizationprocess through deformation in magnetoplastic martensitic materials.

FIGS. 4 (a) and (b) is a schematic depiction of a compressed/indentedcrystal sample.

FIG. 5 is a detail of the shaded face of the sample depicted in FIG. 4(a).

FIGS. 6 (a) and (b) is a schematic depiction of how twin crystallographyis identified. FIG. 6( c) is a graphic depiction of an exampleatomic-force microscopy (AFM) surface profile relating to the anglesφ_(a) and φ_(c) from FIG. 6( b).

FIG. 7 is a schematic depiction of how crystallographic twin variantsare identified.

FIG. 8( a) is a depiction of a magnetic-force microscopy (MFM) result onthe same area displayed in the shaded face (x-z plane) in FIG. 4( aandb) and in FIG. 5. FIG. 8( b) is a magnified view of an area of FIG. 8(a) indicating an interaction of the domains across twin boundaries.

FIG. 9 is a schematic portrayal of one embodiment of the invention,wherein simultaneous writing and reading is performed on the sample ofFIGS. 4 and 5.

FIG. 10 is a table of magnetic and magneto-mechanical properties of somebut not all (potentially) ferromagnetic materials. Values surrounded bya rectangle being least favorable, and values surrounded by a trianglebeing most favorable.

FIG. 11 lists and categorizes many, but not all, magnetoplastic andpotentially magnetoplastic materials, as well as citations to scientificliterature discussing these materials. For materials that are circled inFIG. 11, magnetoplasticity has been demonstrated.

FIG. 12 shows an array of nanoindents performed in the Example below, onthe shaded face of material such as that show by force F in FIG. 4 a.

FIG. 13 (13 a and 13 b) show results from nanoindentation experiments ofthe Example below.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

Referring to the figures, there are shown several, but not the only,embodiments of the invented devices and methods.

The present invention makes use of two effects of materials with mobiletwin boundaries, and, specifically, of magnetoplastic and/ormagnetoelastic materials, including materials from the following subsetsof magnetoplastic and/or magnetoelastic: magnetic shape memory alloys(MSMA, which are ferromagnetic or “magnetic”); materials resulting frommartensitic transformation (typically called “martensite”); and otherproviders of twins.

Magnetoplasticity/magnetoelasticity is the magnetic-field-induceddeformation enabled through the magnetic-field-induced displacement ofcrystallographic twin boundaries. Deformation-induced-change ofmagnetization occurring in magnetoplastic/magnetoelastic materials hasbeen demonstrated by Mullner et al. (Müllner P, Chernenko V A, KostorzG, Scripta Mater 2003b; 49:129). The strain-induced change ofmagnetization due to twin rearrangement is the reverse/inverse effect tomagnetoplasticity. With regard to applications for magnetoplasticity,the present inventors believe that magnetoplasticity is relevant foractuators, whereas the inverse/reverse effect is relevant for sensorsand/or power generation. Numerous publications deal with many aspects ofthe ferromagnetic martensites in Ni—Mn—Ga alloys such as martensitictransformations and martensite structure (e.g. Chernenko et al. 1995),magnetic-field-induced deformation (e.g. Murray et al 2000a, Sozinov etal 2002) and the associated magneto-stress (Chernenko et al. 2000,Mullner et al. 2002, Tickle and James 1999, Chernenko et al. 1999,Murray et al. 2000b). However, only few results concerning theinverse/reverse effect (strain-induced change of magnetization) have sofar been published (Mullner et al 2003b, Soursa et al 2004a, Soursa etal 2004b, Straka et Heczko 2006).

The magnetoplastic effect is related to the magnetic-field-induceddisplacement of twin boundaries, which is a thermodynamicallyirreversible process (Ullakko 1996, Mullner et al. 2002). On themicroscopic scale, a twin boundary moves by the motion of twinningdisconnections (Pond and Celotto 2002), a process which can be triggeredby a magnetic force on the dislocation (Mullner and Ullakko 1998,Ferreira and Vander Sande 1999, Mullner 2006). In Ni₂MnGa, thecooperative motion of twinning dislocations finally leads to a strain ofup to 10% (Müllner et al. 2004a).

FIG. 1 shows the setup of the experiment with which deformation-inducedchange of magnetization was demonstrated. A magnetoplastic material inform of a parallelepiped (1) is fixed between two pressure pistons (2)that introduce the load. Two quartz glass push rods (3) transmit thedisplacement of the top and bottom surfaces of the sample/pistons to theextensometers (not shown on the figure). A Hallbach cylinder(cylindrical permanent magnet, 4) produces the magnetic field H_(x)(large arrow). A Hall probe (5) measures the sum of H_(x) and strayfield H_(S) (small arrows) on the side surface of the sample. Thedifference between measured field and H_(x) is taken as a measure forthe sample magnetization in the direction indicated by the arrows.

Referring to FIG. 1, uniaxial compression experiments under orthogonalmagnetic field were done on a single crystal with compositionNi₅₁Mn₂₈Ga₂₁ (numbers indicate atomic percent). The sample was cut as arectangular prism with {100} faces in all directions and measured5.45(2) mm×3.26(2) mm×2.34(2) mm. In the ferromagnetic austenitic phase,i.e. above the reverse transformation temperature of 316 K, the samplewas a single crystal with the ordered cubic L2₁ structure. At roomtemperature, the material is in the martensitic phase. Thecrystallographic directions a and c of all twin variants were parallelwithin 3° to the sample edges. The easy magnetization axis is parallelto the c direction.

The sample was deformed in uniaxial compression and unloaded at constantspeed (2×10⁻⁶ m/s) in a mechanical testing machine equipped with a 500 Nload cell and extensometers insensitive to magnetic fields. The magneticfield μ₀H=0.7 T was produced by a permanent magnet system. The samplewas mounted in such a way that the longest edge was parallel to thecompression axis (z direction). The x direction was defined parallel tothe shortest edge of the sample; x-y-z constitute Cartesian coordinates.The magnetic field was applied in x direction. A Hall probe waspositioned close to one of the sample surfaces which were parallel tothe y-z plane. The set-up of the experiment is outlined in FIG. 1.

FIG. 2 shows stress (squares) and magnetization (triangles) as afunction of compressive deformation along <100> direction of aNi₅₁Mn₂₈Ga₂₁ (numbers indicate atomic percent) single crystal measuredat constant orthogonal magnetic field of 0.7 T along the x direction.Open and full symbols indicate values for increasing and decreasingdeformation along the z direction. Upon deformation, the stressincreases quickly to about 1.5 MPa at 0.04% strain. Above 0.1% strain,stress increases slowly and almost linearly up to 1.9% strain. At largerstrain, stress increases rapidly again. Upon unloading, the total strainis recovered, however at a lower stress level compared with loading. Themagnetization along the x axis (divided by its value M₀ in theundeformed state) decreases linearly with increasing deformation up to1.9%. Upon unloading, the magnetization restores its initial value witha small hysteresis.

FIG. 2 shows the results of the experimentation represented by FIG. 1.Upon mechanical loading along the z direction at constant speed, thestress increases strongly at the beginning. The slope of thestress-strain curve decreases rapidly right after the onset of plasticdeformation and is almost constant up to about 1.9% compressive strainand a corresponding stress of 6 MPa. At larger strain, the stressincreases more rapidly. Over the whole deformation range, the relativemagnetization in x direction M_(x)/M_(x0)=(H−H_(x))/(H₀−H_(x)) (H and Mare the field detected with the Hall probe and the magnetization of thesample, H₀ and M₀ are the values in the undeformed state) decreaseswithin experimental error linearly with increasing strain. Uponunloading, the stress decreases rapidly at the beginning and more slowlywith decreasing strain until the full deformation is recovered. Therelative magnetization increases again linearly until it reaches theinitial value. The magnetization exhibits a negligible hysteresis. Theslopes of the magnetization curves in both directions of deformation areconstant and equal within experimental error over a wide range of strainwhereas the stress curves have different shapes.

FIG. 3 is a schematic of the magnetization process through deformationin magnetoplastic martensitic materials. (a) Dark and bright grayindicate two twin variants. The local magnetization (arrows) aligns withthe easy axis that is differently oriented for each twin variant. (b)Under an applied magnetic field H_(x), the twin boundaries move causinggrowth of the twin variants with c parallel to the field. In the othertwin variants, the magnetic moments rotate towards the direction of themagnetic field. (c) Under an applied load (F_(z)), the twin boundariesmove causing growth of one twin variant with the crystallographic c axisparallel to the load direction and shrinkage of the other. The specimendeforms since the crystallographic axes c<a. At the same time, thedistribution of magnetic moments changes and alters the totalmagnetization.

The twin rearrangement due to the action of a magnetic field and amechanical force and the associated processes of deformation andmagnetizing on the mesoscopic scale are shown schematically in FIG. 3.In the undeformed state and without magnetic field (FIG. 3 a), the twinstructure contains self-accommodated elastic domains with thecrystallographic c directions distributed irregularly. In the absence of180° magnetic domains (which is always true for magnetic fields of 0.1 Tand larger and even, in some cases, without application of a magneticfield), there is a considerable stray-field H_(S) caused by domains withthe axis of easy magnetization (which is parallel to the c direction)perpendicular to the surface.

In the schematic of FIG. 3, only one martensite domain with one set oftwins is present. Such a structure can be obtained after a suitablemagnetic or mechanical treatment. In the present study, there are manymartensite domains with differently oriented sets of twins.

When a magnetic field H_(x)>>0.1 T is applied along the x direction, thetwin boundaries move in such a way that the twins with c parallel to thex direction grow at the expense of twins with c perpendicular to the xaxis (FIG. 3 b). In regions, through which a twin boundary passes, the cdirection switches from parallel to the z axis to parallel to the xaxis. Since c/a<1, the sample shrinks along the x direction and expandsalong the z direction (magnetoplasticity). In addition to the motion oftwin boundaries, the magnetic field H_(x) causes the magnetic moments inthe domains with c perpendicular to the x direction to rotate by anangle θ with sin θ=H_(x)/H_(A) (H_(A) is the saturation field) towardsthe x direction (FIG. 3 b). In the present experiment, H_(x)/H_(A)≅0.7and θ≅45°. Owing to the rotation of the magnetic moments, thestray-field increases. The Hall probe (FIG. 1) detects the sum of thestray-field and the applied field.

When the sample is mechanically compressed along the z direction, thetwin boundaries move in the opposite direction, i.e. the twins with cparallel to the z direction grow at the expense of twins with cperpendicular to the z axis (FIG. 3 c). In regions, through which a twinboundary has passed, the c direction switched from parallel to the xaxis to parallel to the z axis. Thereby, the direction of the magneticmoments rotates from parallel to the x axis to about 45° inclined to thex axis. This causes a reduction of the stray field which is detected bythe Hall probe. Very close to the sample surface, the magnetic inductionoriginating from the stray field is a linear function of the fractionsof each twin variant. Since the strain is a linear function of the twinvariant fractions, too, the signal of the Hall probe decreases linearlywith strain (FIG. 2). Upon mechanical unloading, the reverse processoccurs. The twin boundaries move again under the action of the magneticfield H_(x) until the twin pattern and the shape of the unloaded state(FIG. 3 b) are reached. Since in the present experiments, strain andmagnetization are controlled by the fractions of twin variants, there isno significant hysteresis between loading and unloading curves (FIG. 2).

The role of the magnetic bias field (H_(x) in the above experiment) istwofold. First, the bias field removes all 180° domain boundaries andcauses a net magnetization M_(x) in x direction. This component of themagnetization induces voltage in the coil, for extraction of electricalpower. Second, the magnetic bias field works against the applied forceand restores the shape of the magnetoplastic and/or magnetoelasticmaterial after removal of the force. The restoration of the shapeimplies a further change of magnetization capable of generatingelectrical power. Below a bias field comparing to the saturation field(about 1 T for Ni₂MnGa), the recoverable strain decreases and vanishesbelow a threshold field (see Mullner et al. 2002). Alternatively, oradditionally, restoration of the initial state may also be achievedthrough the application of a bias stress, for example, by a lever systemforcing the shape of the magnetoplastic and/or magnetoelastic materialback to the same or generally the same shape as the initial state.

FIG. 4 shows a sample geometry and indentation pattern/physical contour.(a) The parallelepiped-shaped sample was compressed along the directionparallel to its shortest edge while heated and cooled. Cartesiancoordinates are defined on the sample as indicated. In most of thevolume, the crystallographic c direction was parallel to the x axis.There are two possibilities for the orientation of the crystallographica and b directions. (b) Schematic of the surface structure withbands/indentation pattern indicated.

A Ni₅₀Mn₂₉Ga₂₁ (numbers indicate atomic percent) single crystal wasgrown following the Bridgman method with a rate of 3.5 mm/h. A samplewas cut using electric discharge erosion forming a parallelepiped withfaces parallel to cubic {001} planes. The composition was measured usingx-ray fluorescence analysis and energy-dispersive x-ray analysis withtransmission and scanning electron microscopes. The composition variedby a couple of percent across the crystal. Electron diffraction revealedthe presence of 14M (orthorhombic) and 10M (tetragonal) martensite withthe 14M martensite being predominant. The dimensions of the sample were5.4 mm×4.0 mm×2.0 mm. Cartesian coordinates are defined on the sample(sample coordinates) such that the shortest, intermediate, and longestedges are parallel to the x, y, and z directions. The crystal wasannealed at 800° C. for one hour under an inert argon atmosphere. Onesurface with dimensions 5.4 mm×2.0 mm was polished with 1 μm diamondslurry (light gray shaded face in FIG. 4). After polishing, the samplewas compressed with forces parallel to the x direction (i.e. parallel tothe shortest edges, see FIG. 4 a) to a stress level of 12±2 MPa andheated to 150° C. During the heating and cooling period, the force washeld constant until the sample reached ambient temperature. The stressof 12 MPa applied during the heat treatment is sufficient to align thecrystallographic c direction with the loading direction (see e.g.Müllner et al. 2004b). Thus, the crystallographic c direction which isthe shortest lattice parameter must be parallel to the x direction inmost of the sample. For the crystallographic a and b directions, thereare two possibilities as indicated in FIG. 4 a.

A Veeco Dimension 3100 Atomic Force Microscopy (AFM) System was used tocharacterize the relief of the surface and the magnetic structure of thesurface. Surface reliefs were taken in tapping mode with a diamond tipwith tip radius 25±3 nm and an apex angle of 60°. For the magnetic-forcemicroscopy (MFM) experiments, a ferromagnetic tip with a Co-coating andmagnetization parallel to the tip axis (i.e. perpendicular to theinvestigated surface) was used. The magnetic tip is particularlysensitive for variations to out-of-plane magnetization.

Long, straight thin bands parallel to the edges were found on thepolished surface (FIG. 4 b). The bands were identified as twins (seebelow). The experiments were terminated when an accidental collision ofthe AFM with the sample surface occurred. After the collision, the twinboundaries have moved and the twins have thickened considerably.

After the thermo-mechanical treatment, the polished surface displayed analmost periodic pattern of thin criss-crossing bands parallel to thesample edges. FIG. 4 b is a schematic of the polished surface with“long” twins (i.e. twins parallel to the long edge of the sample whichis the z direction) shaded gray and “short” twins (i.e. twins parallelto the short edge of the sample) shaded black.

FIG. 5 is a section of the surface relief of the shaded face in FIG. 4a, wherein the image is produced by atomic force microscopy (AFM). Thesurface relief contains terraces separated by slopes stemming fromtwins, which are the bands in FIG. 4 b. FIG. 5 is a surface relief takenwith the AFM in tapping mode. The surface shows terraces separated byinclined slopes. The root mean square surface roughness is 8 nm.

FIG. 6 shows how the twin crystallography is identified. (a)Cross-section of a twin parallel to the z direction. The orientation ofeach twin variant is indicated. (b) Cross-section of a twin parallel tothe x direction. (c) Example of a AFM surface profile used forevaluating the relief angles φ_(a) and φ_(c). The slope angles φ_(a) andφ_(c) (defined in FIGS. 6 a,b) were evaluated from surface profiles(e.g. FIG. 6 c) and are (φ_(a)=3.0°±0.3° and φ_(c)=4.3°±0.5°. The slopeangles are related to the lattice parameters byφ_(a)=arctan b/c−arctan c/b  (6a)φ_(c)=arctan a/b−arctan b/a  (6b)where the orientation of the crystallographic directions is given inFIGS. 6 a and b. There are, in principal, many possibilities for thearrangement of twins and the orientation of the crystallographicdirections in each twin. However, the crystallographic c direction wasoriented by applying a normal stress of 12 MPa along the x direction(FIG. 4 a). Thus, the predominant lattice parameter parallel to the xdirection in the crystal must be the c lattice parameter which is theshortest.

FIG. 7 shows the crystallography and nomenclature of twin variants.Orientation and gray scale corresponds to FIG. 1 b. Lower case a, b andc indicate the orientation of the crystallographic directions. Capitalletters mark the crystallographic direction perpendicular to the surfaceand give the variants the name. Since the slope angle for an ac twin(with the b direction parallel to the twin boundary) is more than 6°(assuming c/a=0.9), the presence of ac twins is excluded. The onlycombination of crystallographic orientations is a∥z, b∥x, and c∥y (Bvariant, since b is perpendicular to the surface) in the largestregions, a∥z, b∥y, and c∥x (C variant) in the long twins, and a∥x, b∥z,and c∥y (A variant) in the short twins (FIG. 7).

FIG. 8 shows results of magnetic-force microscopy (MFM) on the same areaas displayed as relief in FIG. 5. FIG. 8( a) is a MFM image of the areashown in FIG. 5. The C variants appear as dark lines indicating stronginteraction with the MFM tip and out-of-plane magnetization. The A and Bvariants are hardly distinguishable. Domains are elongated along thecrystallographic c direction. FIG. 8( b) is a magnified view of an areasimilar to FIG. 8( a) indicating an interaction of the domains acrosstwin boundaries.

FIG. 8 a is an MFM image of the same area as shown in FIG. 5. The strongcontrast in the two horizontal stripes along the z direction indicatesout-of-plane magnetization in the long twins. For these twins, thecrystallographic c direction is perpendicular to the surface (Cvariants), i.e. parallel to the x direction (compare FIG. 7). Thecrystallographic c direction is the direction of easy magnetization(Sozinov and Ullakko 2002). The regions with the A and B variants (FIG.5) display a weak magnetic contrast in form of long stripes. FIG. 8 b isa magnified view of FIG. 8 a also containing regions with A, B and Cvariants. FIG. 8 b was taken with higher magnetic sensitivity comparedto FIG. 8 a. The stripes are parallel to the y direction which isparallel to the crystallographic c direction in both A and B variants(FIG. 7).

For a cubic-to-orthorhombic martensite transformation, there are sixorientation variants in the martensitic phase. Each variant has six{110}<1 10> twinning systems; two of each ab-, bc-, and ac-twin type(Müllner et al. 2004b, Müllner et al. 2005). Thus, there are 36 possibletwinning systems not including twins forming within twin variants. It isremarkable that given one known crystallographic direction in thepredominant martensite variant, which in the present case is c, theanalysis of the surface profile yields the complete crystallography ofall twin variants. Thus, crystallographic orientation and spatialconfiguration of the twin structure can be fully characterized by AFMalone without employing diffraction methods.

The identification of the A, B, and C variants and the orientation ofthe twin boundaries prove the presence of ab-twins and bc-twins. Forab-twins (bc-twins), the twin boundary contains the crystallographic c(a) direction. It may be noted that the stress induced by the AFMaccidentally impacting the sample (as described above) was sufficient tomove the twin boundaries. Thus, ab-twins and bc-twins are very mobile.The present results substantiate the idea that different twinning modesare active in orthorhombic Ni—Mn—Ga MSMA resulting in a variation of themagnetic-field-induced strain.

The magnetic-force microscopy images reveal a strong signal for the Cvariants. Since the magnetization of the MFM tip is perpendicular to thesurface, a strong signal indicates an out-of-plane magnetization, whichis parallel to the crystallographic c direction. The c direction is thedirection of easy magnetization in orthorhombic Ni—Mn—Ga martensite(Sozinov and Ullakko 2002). Thus, the conclusion that the C variant hasout-of-plane magnetization is consistent with the magnetic properties ofthe material. The magnetic contrast in A and B variants reveals magneticdomains which are elongated along the crystallographic c direction. Thewidth of the magnetic domains is the same for A and B variants. Thisindicates that magnetic domains have a rod-like shape. The domains matchacross twin boundaries (FIG. 6 b) to ensure magnetic compatibility byavoiding magnetic charges on the twin boundaries (Vlasova et al. 2000).Similar magnetic domain structures were reported for tetragonal Ni—Mn—Gamartensite (Ge et al. 2004, Pan and James 2000, Sullivan and Chopra2004).

Thus, it has been shown above on macroscopic samples that it is possibleto control the magnetic and crystallographic states of a magnetoplasticmaterial through the application of magnetic fields and/or mechanicalforces. This represents one embodiment of a writing process. It has alsobeen shown that it is possible to read the state of individual twindomains through the use of atomic force microscopy and magnetic forcemicroscopy. A schematic of apparatus and methods of writing on (W, herea mechanical force F) and reading of (R, here apparatus for atomic forceor magnetic force microscopy) a macroscopic portion of multi-statematerial is shown in FIG. 9. While FIG. 9 portrays a mechanical forcefor W, it will be understood from this disclosure and from FIG. 1 that amagnetic-field-change system may also be used for the writing system W.

With current micro-fabrication tools, micro- and nano-pillars withdiameters of 100 nm and below can be produced. For instance, goldpillars with thickness between 100 nm and 200 nm were recently produced,mechanically deformed using a nanoindenter and characterized (Greer etal. 2005). Thus, the writing method outlined above and depicted in FIGS.1 and 2 may directly be downscaled to the nano- and micro-scales.

One embodiment of the system uses mechanical force with or withoutcombined use of magnetic field to write the “bits” on/in magnetoplasticand/or magnetoelastic material. We do not use “bit” in this disclosureto refer to the traditional meaning referring to binary coding, but,instead, we use it for convenience to refer to the multi-state elementsof embodiments of the present invention, wherein said multi-statespreferably comprise greater than two states, and more preferably six ormore for a single magnetoplastic/magnetoelastic element. Likewise, theterm “element” is not intended to refer to elements of the PeriodicTable, but rather individual pieces/portions/regions (for example,individual memory units) according to the invention that are preferablyformed/defined by pieces/portions/regions ofmagnetoplastic/magnetoelastic material.

One embodiment of the system uses a reading process that may be achievedthrough a combination of atomic-force microscopy and magnetic-forcemicroscopy; diffraction-based work/reading is not required in preferredembodiments. In other embodiments, reading is accomplished through othermeans including optical interferometry, spin-polarized scanning electronmicroscopy, the magneto-optical Kerr effect, and/or combination of suchmethods.

One example, of how multi-functionality of the embodiments of theinvention may be used, is as follows: one functionality may comprise theshape state serving as the main actuation function such as positioning aprobe or defining a memory. Simultaneously, another functionality may beprovided, for example, the magnetic state providing a signal to indicatethe current position. In this embodiment, displace function and indicatefunction are combined in one material. Other embodiments combine otheroperations such as read & write, sense & indicate, sense & control, forexample.

Preferred Materials:

The materials used for the preferred multi-state memory andmulti-functional devices are those which couple magnetic andcrystallographic/shape states. The preferred materials are those withmobile twin boundaries, wherein strain-induced deformation/change insaid twin boundaries produces change in magnetization, and whereinmagnetic-field change produces deformation/change in the twin boundaries(the later being the feature that is typically implied by the terms“magnetoplasticity” or “magnetoelasticity.”). Therefore, the preferredmaterials used for the multi-state memory and/or multi-functionaldevices of the invention may be selected from the broad categories ofmagnetoplastic and magnetoelastic materials, including from thefollowing subsets of magnetoplastic and/or magnetoelastic: magneticshape memory alloys (MSMA, which are ferromagnetic or “magnetic”);materials resulting from martensitic transformation (typically called“martensite”); and other providers of twins. The terms “magnetoplastic”and “magnetoelastic” overlap to some extent, in that “magnetoelastic” iscommonly used even for some materials that do exhibit hysteresis.

The phrases “magnetplasticity/magnetoelasticity” and“magnetoplatic/magnetoelastic materials” as applied in embodiments ofthe invention do not necessarily include all ferromagnetic materials,for example, these phrases preferably do not include “classical”magnetostriction (magnetostricitive materials). Whereas currentmagnetostrictive materials are limited to a maximum of about 0.2% strain(for Terfenol-D), the preferred megnetoplastic/megnetoelastic materialsexhibit typically above 1% strain, and, in some embodiments, up to about10% strain and possibly more.

Further, the magnetoplastic/magnetoelastic materials used in preferredembodiments of this invention are not piezoelectric. However, theinventors envision that there may be materials developed or discoveredin the future that are both magnetoplastic/magnetoelastic andpiezoelectric, and, hence, could be included in the preferredembodiments. The inventors envision that materials may be recognized ordiscovered in the future that are magnetoplastic/magnetoelastic andpiezoelectric, and that may offer more memory states (electrical states)and also more actuation options (namely electrical field) for switchingfrom one state to another (control).

Requisite for magnetoplasticity is the magnetic-field-induced motion oftwin boundaries. This requisite implies the following properties:

-   (i) The material must deform more easily by twinning than by    dislocation motion.-   (ii) The twinning stress must be less than the magnetostress τ_(M),    i.e. the stress which can be induced through a magnetic field.    Factors affecting (i) above include the crystal structure (trend:    lower symmetry is better than higher symmetry), the lattice    potential (trend: strong bonding is better than weak bonding), the    size of the lattice parameter (trend: larger is better than    smaller). Factors affecting (ii) include the magnetic anisotropy    constant K (the higher the better) and the twinning shear (the    smaller the better).

Regarding applications, a large strain might be desirable. This impliesa large twinning shear, which is in conflict with (ii). Furthermore, itis desirable to obtain magnetoplasticity with a small magnetic field.This implies a large saturation magnetization M_(s). Thus, the desiredmaterials properties are:

-   -   1. Large magnetic anisotropy K.    -   2. Large saturation magnetization M_(s).    -   3. For large stress output: small twinning shear.    -   4. For large strain output: large twinning shear.

Regarding memory applications of preferred embodiments of the invention,a large strain output is important; however, the requirement that thestress required for twinning must be smaller than the magnetostressstill holds. The multi-state data storage medium according to preferredembodiments of the invention should be material with the basicproperties (i), (ii), and (1-4) above, and may be manufactured as solidbulk part, as thin film, as nanostructured bulk part, as nanostructuredthin film, as nano-columns and in any other shape and form.Nano-structured arrays of the material may be of particular interest asthe bit size can be tailored in this case using nano-fabricationmethods.

FIG. 10 summarizes magnetic and magneto-mechanical properties of some(potentially) ferromagnetic materials. Materials in FIG. 10 for whichmagnetoplasticity has been reported in the literature are circled.Embodiments of the invention may include one or more of the materials inFIG. 10, with values surrounded by a rectangle being least favorable,and values surrounded by a triangle being most favorable. The strainε_(M,max) is proportional to the twinning shear and marks thetheoretical maximum of magnetic-field-induced strain. The saturationfield μ₀H_(a) is the magnetic field at which the maximum magnetostressτ_(M,max) is reached. Further increase of the magnetic field does notincrease the magnetostress.

Current research in the field of magnetoplasticity focuses onferromagnetic shape-memory alloys (MSMA), because in these materials,the twinning stress is very low. Particular attention is being paid toHeusler alloys, particularly off-stoichiometric Ni₂MnGa. Otherferromagnetic shape-memory alloys (i.e. non-Heusler alloys) which areunder study include FePd, CoPt, FePt, and Fe₃Pd. Recently,magnetoplasticity was reported for an antiferromagnetic (AFM) magneticshape-memory alloy γ-Mn—Fe—Cu (J. H. Zhang, W. Y. Peng, S. Chen, T. Y.Hsu (X. Zaoyao), Appl. Phys. Lett. 86, 022506 (2005)). Non-shape-memoryalloys which have been studied in context of magnetoplasticity includedysprosium and τ-MnAl—C.

FIG. 11 lists and categorizes many magnetoplastic and potentiallymagnetoplastic materials, as well as citations to scientific literaturediscussing these materials. Embodiments of the invention may include oneor more of the listed materials and/or one or more materials selectedfrom the broad categories of materials. For materials that are circledin FIG. 11, magnetoplasticity has been demonstrated. The citations inFIG. 11 are:

-   [Cui 2004] J. Cui, T. W. Shield, R. D. James, Acta mater. 52, 35    (2004).-   [Fujita 2000] A. Fujita, K. Fukamichi, F. Gejima, R. KainLima, K.    Ishida, Appl. Phys. Lett. 77, 3054 (2000).-   [James 1998] R. D. James and M. Wuttig, Phil. Mag. A 77, 1273    (1998).-   [Kostorz 2005] G. Kostorz and P. Müllner, Z. f. Metallk. 96, 703    (2005).-   [Lieb. 1976] H. H. Lieberrnann and C. D. Graham, Jr. Acta Met. 25,    715 (1976).-   [Santa. 2006] R. Santamarta, E. Cesari, J. Font, J. Muntasell, J.    Pons, J. Dutkiewicz, Scripta Mater. 54, 1985 (2006).-   [Solo. 2004] A. S. Sologubenko, P. Müllner, H. Heinrich, K.    Kostorz, Z. f. Metallk. 95, 486 (2004).-   [Vlasova 2000] N. I. Vlasova, G. S. Kandaurova, N. N.    Shchegoleva, J. Magn. Magn. Mater. 222, 138 (2000).-   [Wada 2003] T. Wada, T. Tagawa, M. Taya, Scripta Mater. 48, 207    (2003).-   [Wuttig 2001] M. Wuttig, J. Li, C. Craciunescu, Scripta Mater. 44,    2393 (2001).-   [Zhang 2005] J. H. Zhang, W. Y. Peng, S. Chen, T. Y. Hsu (X.    Zaoyao), Appl. Phys. Lett. 86, 022506 (2005).

As may be seen from the above comments and FIGS. 10 and 11, preferredmaterials may be selected from the group consisting of: Heuslershape-memory alloy, non-Heusler ferromagnetic shape-memory alloy,antiferromagnetic shape-memory alloy, non-shape-memory magnetoplasticalloy, Ni₂MnGa, Dy, α-Fe, Co—Ni, τ-MnAl—C, L1₀ FePd, L1₀CoPt, L1₀FePt,Ni₂MnGa, Co₂NiGa, Ni₂MnAl, Ni₂FeGa, Fe₃Pd, Fe—Pd—PT, γ-Mn—Fe—Cu,Ni₅₁Mn₂₈Ga₂₁, Ni—Mn—Ga alloys, Ni—Mn—Ga ferromagnetic martensite, andcombinations thereof. Especially-preferred for non-volatile memoryembodiments will be materials, including materials selected from thebroad category of magnetoplastic/magnetoelastic and from the above list,that exhibit magnetic and crystallographic states that are stablewithout the use of an external electrical input or other force.

DISCUSSION OF FEATURES AND ADVANTAGES OF PREFERRED EMBODIMENTS

Current conventional device, logical and memory technology is based onbinary elements. On the other hand, elements, according to embodimentsof the invention, made of magnetic shape-memory material have sixcrystallographic states. The crystallographic states are linked to threemagnetic states. This enables the combined usage of a single element asmemory and logic unit. As a consequence, the low-level programming ofcomputers becomes much more efficient. Because a single element may havesix states, memory density increases 3 fold.

Current memory devices including dynamic random access memory (DRAM) arebased on volatile mechanisms. Those mechanisms depend on the permanentsupply of power. When power fails, data is lost. The multi-state memoryof preferred embodiments of the invention may be based on materials,from the broad category of magnetoplastic/magnetoelastic materials,selected for exhibiting magnetic and crystallographic states that arestable without the use of an external force (e.g., external electricfield); such multi-state memory may, therefore, be non-volatile

Embodiments of the present invention may continue the historic trendtoward functionalities being carried by materials rather than by machineparts. Further, embodiments of the invention, by providing multiplefunctions from one material, may reduce the complexity of systems whichmay lead to an increase of performance, miniaturization, and lowering ofcosts.

EXAMPLE

The following is an example comprising study and analysis of material,and comprising writing and reading methods of the multi-states of saidmaterial, that may be applied to one, but not the only, embodiment ofthe invention.

Atomic-force microscopy, magnetic-force microscopy (MFM), andnanoindentation experiments were performed on a Ni—Mn—Ga single crystalwith orthorhombic (14M) martensite. The surface relief was characterizedand used to identify the orientation of individual twin variants. It wasshown that ab- and bc-twinning modes are present. The magnetic domainsare elongated along the crystallographic c direction. Twin variants withthe c direction perpendicular to the surface exhibit an out-of-planemagnetization which can be identified through a strong contrast in theMFM image. The deformation caused with nanoindentation is recoverable toa large degree. No trace of deformation twinning was found afterunloading. The results are discussed in the light of pseudo-elastictwinning and dislocation activity. It is concluded that although theresidual deformation is due to dislocation activity, it is likely thatpseudo-elastic twinning took place during loading and unloading. Theresults indicate a size effect on deformation twinning under localizedloading.

1. Introduction

Magnetic shape-memory alloys (MSMA) are distinguished from otherferromagnetic materials by the coupling mechanism between shape andmagnetization. Large (uniaxial) magnetic anisotropy constant and mobiletwin boundaries are at the origin of large magnetic-field-induceddeformation (magnetoplasticity). Plastic (i.e. irreversible) strains upto 10% have been reported for off-stoichiometric Ni₂MnGa Heusler alloysingle crystals (see e.g. references [1-4], below). The reverse effector deformation-induced change of magnetization requires a magnetic fieldbias to minimize demagnetizing by the formation of 180° domains(references [5-7]).

There are various ideas for applications of MSMA. For instance, actuatorapplications including stepper motors, pumps, valves and devices forsurgery make use of the “direct” effect of magnetic-field-induceddeformation. Sensor applications make use of the “reverse” effect ordeformation-induced change of magnetization. In all application,considerable magnetic fields need to be produced. Devices involving MSMAtransducers in the millimeter size range and larger may requireconsiderable power. In addition, losses from induction limit theactuation frequency of large systems. That these obstructions arereduced when the scale is smaller provides motivation forminiaturization. However, mechanical properties change dramatically whenobjects are small (at least in one direction) compared to acharacteristic length specific for a certain deformation mechanism[reference 8], and the yield stress may increase strongly withdecreasing size [reference 9]. Thus the investigation of mechanicalproperties on the micrometer and smaller scales is needed for thedevelopment of miniaturized applications.

The micromechanism of magnetoplasticity is the motion of twinningdisconnections [references 10-14]. Disconnections are interfacial linedefects with elastic fields (displacement, strain and stress) identicalto the elastic fields of dislocations. Greer et al. [reference 9] assumethat the efficiency of dislocation multiplication decreases withdecreasing size of single crystalline gold columns with diameter below 1μm. Lack of dislocations is the reason for the high yield strength ofthese gold columns when only a few hundred nanometer in diameter.

Whether such observations apply to twinning in magnetic shape-memoryalloys and what limits magnetic-field-induced twinning anddeformation-induced change of magnetization at small length scale areopen questions. The present paper reports on an investigation of twoquestions, (i) whether it is possible to locally induce deformationtwinning with a sharp tip thereby inducing a magnetic signature whichcan be identified with magnetic-force microscopy (MFM) and (ii) if thereis a size effect on the deformation mechanisms which are triggered dueto localized loading.

2. Magnetoplasticity

2.1 Micromechanism of Magnetoplasticity

The magnetoplastic effect is related to the magnetic-field-induceddisplacement of twin boundaries. On the microscopic scale, a twinboundary moves by the motion of twinning disconnections [references10-16], a process which can be triggered by a force on thedisconnection. A disconnection is an interfacial line defect with adislocation component and a step component. A disconnection ischaracterized by a line vector (which is parallel to the interface), astep vector (which is perpendicular to the interface) and a Burgersvector (which is of arbitrary direction). For a twinning disconnection,the Burgers vector is parallel to the interface. The motion of atwinning disconnection is conservative and compares to glide of alattice dislocation.

The Burgers vector defines the deformation carried by a movingdisconnection. Thus, the dislocation component controls the interactionof a stress field with the disconnection. A shear stress τ exerts amechanical force F_(mech) on a disconnection:F_(mech)=τb  (1)where b is the magnitude of the Burgers vector. The step vector definesthe volume which is transformed by a moving disconnection. Thus, thestep component controls the interaction of a magnetic field with thedisconnection. For a material with uniaxial magnetic anisotropy K, themagnetic force F_(mag) exerted by a magnetic field on the disconnectionwas derived in [reference 12]. In the special case where the magneticfield is parallel and perpendicular to the easy axis of magnetization ofthe twinned crystals sharing the interface, the magnetic force is

$\begin{matrix}{F_{mech} = \left\{ \begin{matrix}{\mu_{0}{{MHd}\left( {1 - \frac{\mu_{0}{MH}}{4\; k}} \right)}} & {{{for}\mspace{14mu} H} \leq H_{A}} \\{Kd} & {{{for}\mspace{14mu} H} \geq H_{A}}\end{matrix} \right.} & (2)\end{matrix}$where μ₀ is the free space permeability, M is the saturationmagnetization, H is the magnetic field, d is the step height of thedisconnection, and H_(A)=2K/μ₀M is the saturation field. Equatingmechanical and magnetic force yields the magnetostress which is themechanical stress that is produced by a magnetic field. The maximummagnetostress is the ratio K/s where s=b/d is the twinning shear.2.2 Twin-Surface Interaction

When a twin grows from a surface into a crystal, the energy increases asa result of three factors: the growing interface, the growing strainfield and, if the material is ferromagnetic, the demagnetizing field.The energy contribution of the twin boundary is proportional to theinterface area whereas the work done by the growing twin is proportionalto the twinned volume, the twinning shear and the applied stress.Therefore, the shear stress τ required to grow the twin decays withincreasing twin thickness t:

$\begin{matrix}{\tau = \frac{\alpha\;\gamma_{tb}}{st}} & (3)\end{matrix}$where α is a geometrical factor on the order of utility, γ_(tb) is thetwin boundary energy, and s is the twinning shear.

The disconnections experience an image force due to the dislocationcomponent. The image force is inversely proportional to the distancefrom the surface and tends to pull dislocations out of the crystal (e.g.reference [17]). When a deformation twin is formed from a surface andgrows into the crystal, the twin disappears after unloading if the imageforce on each dislocation is larger than the force due to internalstresses. This effect may be called elastic twinning [reference 18] orpseudo-elastic twinning. Pseudo-elastic twinning was observed by Garberin calcite in the late 1930s. Garber published his observations inRussian. Garber's results were reviewed in English by Kosevich[reference 18]. Pseudo-elastic twinning requires stress concentrationson the surface [reference 18].

The energy contribution due to the reorientation of the axis of easymagnetization depends on the magnetic domain structure and the relativesizes of twin and magnetic domains. If a magnetic field larger than thesaturation field is applied, the magnetic effects on twin growth arelimited to the magnetic force on twinning disconnections, Eq. (2). Thesame is true for a moderate magnetic field as long as the field is largeenough to remove 180° domain boundaries. For weak magnetic fields, thereorientation of the axis of easy magnetization may result in localizedmagnetic domains with out-of-plane magnetization.

These domains have a strong magnetic field with large energy. Thisenergy provides a further driving force for detwinning after unloadingand may contribute to pseudo-elastic twinning.

2.3 Twin-Interface Interaction

Interfaces interrupt the slip plane (which for twinning disconnectionsis the twinning plane) and, thus, determine the free path of twinningdisconnections. Interfaces themselves do not have long-range stressfield. Therefore, a ‘first’ twinning disconnection that approaches a‘defect free’ interface (i.e. an interface that does not already containa twinning disconnection) moves under the action of a magnetic ormechanical force until it is blocked by the interface. The stress fieldof the blocked dislocation obstructs the motion of the second twinningdisconnection. A critical mechanical or magnetic stress τ_(c) isrequired, to push the second disconnection into the interface. Thecritical stress is inversely proportional to the total number N=d/t oftwinning disconnections [references 19, 20]:

$\begin{matrix}{\tau_{c} = \frac{Gd}{4\sqrt{2}\left( {1 - v} \right)t}} & (4)\end{matrix}$where G and v are the shear modulus and Poisson's ratio, d and t are thed-spacing of twining planes and the twin thickness. For very small twins(e.g. t=10 nm) the critical shear stress τ_(c) is about 10⁻² G, which istwo orders of magnitude larger than the magnetostress. Thus, the seconddisconnection of a thin twin does not reach the interface. For verythick twins (e.g. t=100 μm) the critical shear stress τ_(c) is about10⁻⁶ G, which is a factor of 30 less than the magnetostress. Thus, thesecond, third and following disconnections reach the interface and forma nascent dislocation wall. As the dislocation wall grows, the moredisconnection that are pushed into the interface, the stronger is therepulsive force that these disconnection exert on a incomingdisconnections. The fraction η of twinning disconnections that form adisconnection wall in an interface can be approximated as [references14, 20, 21]:

$\begin{matrix}{\eta = {1 - {\sqrt{2}s\frac{\tau_{c}}{\tau}}}} & (5)\end{matrix}$where τ(>τ_(c)) is the resolved shear stress on the twinning plane inthe twinning direction.

A dislocation wall is a mechanically stable dislocation configuration.Dislocations are trapped in a dislocation wall. It therefore requires aforce to remove a dislocation from a wall. This force increases withincreasing wall size and causes hysteresis [reference 13].

3. Experiments

A Ni₅₀Mn₂₉Ga₂₁ (numbers indicate at. %) single crystal was grownfollowing the Bridgman method with a rate of 3.5 mm/h. A sample was cutusing electric discharge erosion forming a parallelepiped with facesparallel to cubic {001} planes. The composition was measured using x-rayfluorescence analysis and energy-dispersive x-ray analysis withtransmission and scanning electron microscopes. The composition variedby a couple of percent across the crystal. Electron diffraction revealedthe presence of 14M (orthorhombic) and 10M (tetragonal) martensite withthe 14M martensite being predominant. The dimensions of the sample were5.4 mm×4.0 mm×2.0 mm. Cartesian coordinates are defined on the sample(sample coordinates) such that the shortest, intermediate, and longestedges are parallel to the x, y, and z directions. The crystal wasannealed at 800° C. for one hour under an inert argon atmosphere. Onesurface with dimensions 5.4 mm×2.0 mm was polished with 1 μm diamondslurry (light gray shaded face in FIG. 4 a). After polishing, the samplewas compressed with forces parallel to the x direction (i.e. parallel tothe shortest edges, see FIG. 4 a) to a stress level of 12±2 MPa andheated to 150° C. During the heating and cooling period, the force washeld constant until the sample reached ambient temperature. The stressof 12 MPa applied during the heat treatment is sufficient to align thecrystallographic c direction with the loading direction (see e.g.reference [4]). Thus, the crystallographic c direction which is theshortest lattice parameter must be parallel to the x direction in mostof the sample. For the crystallographic a and b directions, there aretwo possibilities as indicated in FIG. 4 a.

FIGS. 4 a and b: Sample geometry and indentation pattern. (a) Theparallelepiped-shaped sample was compressed along the direction parallelto its shortest edge while heated and cooled. Cartesian coordinates aredefined on the sample as indicated. In most of the volume, thecrystallographic c direction was parallel to the x-axis. There are twopossibilities for the orientation of the crystallographic a and bdirections. (b) Schematic of the surface structure with bands.

A Veeco Dimension 3100 Atomic Force Microscopy (AFM) System was used tocharacterize the relief and the magnetic structure of the surface and toperform nanoindentation experiments. Surface relief were taken intapping mode with a diamond tip with tip radius 25±3 nm and an apexangle of 60°. For the nanoindentation experiments, diamond tip,cantilever and photo detector were calibrated using a sapphire singlecrystal. The tip was pressed onto the sapphire surface up to a loadcorresponding to a photo detector reading of 2 V. Up to this load, thediamond tip did not indent the sapphire surface. For the magnetic-forcemicroscopy (MFM) experiments, a ferromagnetic tip with a Co-coating andmagnetization parallel to the tip axis (i.e. perpendicular to theinvestigated surface) was used. The magnetic tip is particularlysensitive for variations to out-of-plane magnetization.

Long, straight thin bands parallel to the edges were found on thepolished surface (FIG. 4 b). In the regions bound by the thin bands andwithin the bands themselves, arrays of nanoindentation experiments wereperformed. FIG. 12 shows an array of indents. Each triangular shapedindent is surrounded by pileups of displaced or deformed material. Ineach line of 10 experiments, the maximum applied load increase fromlight to left from a photo detector reading of 0.2 V (corresponding to6.1 μN) in increments of 0.2 V to a maximum reading of 2.0 V(corresponding to 61 μN). The bands were identified as twins (seeparagraph 4). Indents were set at least 1 μm apart from each other andalso from the nearest twin boundary. In a second series, the distancefrom the nearest twin boundary was varied from indenting on the boundaryto a distance of more than 1 μm. Magnetic-force microscopy images weretaken of some of the indents which resulted from loading up to a photodetector reading of 2 V.

FIG. 12: Array of nanoindents performed on the shaded face in FIG. 4 a.The indents to the outermost right (barely visible) were obtained withloading to a maximum detector voltage of 0.2 V corresponding to 6.1 μN.The maximum load increases from indent to indent towards the left by 6.1μN and measures 61 μN for the indents on the outermost left.

The experiments were terminated when an accidental collision of the AFMwith the sample surface occurred. After the collision, the twinboundaries have moved and the twins have thickened considerably.

4. Results

After the thermo-mechanical treatment, the polished surface displayed analmost periodic pattern of thin crisscrossing bands parallel to thesample edges. FIG. 4 b is a schematic of the polished surface with“long” twins (i.e. twins parallel to the long edge of the sample whichis the z direction) shaded black and “short” twins (i.e. twins parallelto the short edge of the sample) shaded gray. FIG. 5 is a surface relieftaken with the AFM in tapping mode. The surface shows terraces separatedby inclined slopes. The root mean square surface roughness is 8 nm. Theslope angles φ_(a) and φ_(c) (defined in FIGS. 6 a and b) were evaluatedfrom surface profiles (e.g. FIG. 6 c) and are φ_(a)=3.0°±0.3° andφ_(c)=4.3°±0.5°. The slope angles are related to the lattice parametersby

$\begin{matrix}{\varphi_{a} = {{\arctan\frac{b}{c}} - {\arctan\frac{c}{b}}}} & \left( {6\; a} \right) \\{\varphi_{c} = {{\arctan\frac{a}{b}} - {\arctan\frac{b}{a}}}} & \left( {6\; b} \right)\end{matrix}$where the orientation of the crystallographic directions is given inFIGS. 6 a, b.

FIG. 5: Surface relief of the shaded face in FIG. 4 a. The surfacerelief contains terraces separated by slopes stemming from twins, whichare the bands in FIG. 4 b.

FIG. 6: Schematic of twin crystallography. (a) Cross-section of a twinparallel to the z direction. The orientation of each twin variant isindicated. (b) Cross-section of a twin parallel to the x direction. (c)Example of a surface profile used for evaluating the relief angles φ_(a)and φ_(c).

There are, in principal, many possibilities for the arrangement of twinsand the orientation of the crystallographic directions in each twin.However, the crystallographic c direction was oriented by applying anormal stress of 12 MPa along the x direction (FIG. 4 a). Thus, thepredominant lattice parameter parallel to the x direction in the crystalmust be the c lattice parameter which is the shortest. Since the slopeangle for an ac twin (with the b direction parallel to the twinboundary) is more than 6° (assuming c/a=0.9), the presence of ac twinsis excluded. The only combination of crystallographic orientations isa∥z, b∥y, and c∥x (B variant, since b is perpendicular to the surface)in the largest regions, a∥z, b∥x, and c∥y (C variant) in the long twins,and a∥y, b∥z, and c∥x (A variant) in the short twins (FIG. 7).

FIG. 8 a is an MFM image of the same area as shown in FIG. 5. The strongcontrast in the two horizontal stripes along the z direction indicatesout-of-plane magnetization in the long twins. For these twins, thecrystallographic c direction is perpendicular to the surface (Cvariants), i.e. parallel to the y direction (compare FIG. 7). Thecrystallographic c direction is the direction of easy magnetization[reference 3]. The regions with the A and B variants (FIG. 7) display aweak magnetic contrast in form of long stripes. FIG. 8 b is a magnifiedview of FIG. 8 a also containing regions with A, B and C variants. FIG.8 b was taken with higher magnetic sensitivity compared to FIG. 8 a. Thestripes are parallel to the y direction which is parallel to thecrystallographic c direction in both A and B variants (FIG. 7).

FIG. 7: Crystallography and nomenclature of twin variants. Orientationand gray scale corresponds to FIG. 4 b. Lower case a, b, and c indicatethe orientation of the crystallographic directions. Capital letters markthe crystallographic direction perpendicular to the surface and give thevariants the name.

FIGS. 8 a and b: Magnetic-force microscopy. (a) MFM image of the areashown in FIG. 5. The C variants appear as dark lines indicating stronginteraction with the MFM tip and out-of-plane magnetization. The A and Bvariants are hardly distinguishable. Domains are elongated along thecrystallographic c direction. (b) Magnified view of an area similar to(a) indicating an interaction of the domains across twin boundaries.

FIG. 13 a displays the results of the nanoindentation experiments.Triangles, squares and circles refer to experiments in A, B, and Cvariants. Open symbols indicate the indentation depth at full loadwhereas frill symbols indicate the residual indentation depth afterunloading. The indentation depth at full load increases almost linearlyfor all variants. The depth is deepest in the B variants and least deepin the A variants. The same trend is observed for the depth of theresidual indent. However, the depth of the residual indent amounts toonly about 30% of the depth at maximum load. The results of indentationexperiments at different distances from the twin boundaries wereidentical within experimental error. FIG. 13 b is the stress versusindentation depth diagram for the B domain. The stress represents anormalized value obtained from dividing the applied load by the area ofthe indent. The stress is initially very large and decreases withincreasing depth.

FIGS. 13A and B: Results from nanoindentation experiments. (a)Indentation depth vs. maximum load. The open symbols represent the depthat maximum load; the full symbols mark the residual depth afterunloading. Triangles, squares and circles indicate results from A, B,and C variants. The depth at maximum load increases nearly linearly withmaximum applied load. The indentation depth in C variants issystematically larger than the depth in B and A variants. The residualdeformation is smallest for A variants and about the same for B and Cvariants. (b) Data of B variants evaluated for stress vs. indentationdepth. With increasing depth the average stress under the indenterdecreases.

5. Discussion

For a cubic-to-orthorhombic martensite transformation, there are sixorientation variants in the martensitic phase. Each variant has six{110}<1 10> twinning systems; two of each ab-, bc-, and ac-twin type[references 4, 13]. Thus, there are 36 possible twinning systems notincluding twins forming within twill variants. It is remarkable thatgiven one known crystallographic direction in the predominant martensitevariant, which in the present case is c, the analysis of the surfaceprofile yields the complete crystallography of all twin variants. Thus,crystallographic orientation and spatial configuration of the twinstructure can be fully characterized by AFM alone without employingdiffraction methods.

The identification of the A, B, and C variants and the orientation ofthe twin boundaries prove the presence of ab-twins and bc-twins. Forab-twins (bc-twins), the twin boundary contains the crystallographic c(a) direction. The stress induced by the AFM accidentally impacting thesample was sufficient to move the twin boundaries. Thus, ab-twins andbc-twins are very mobile. It was suggested in [reference 4] thatdifferent twinning modes are active in orthorhombic Ni—Mn—Ga MSMAresulting in a variation of the magnetic-field-induced strain. Thepresent results substantiate that suggestion.

The magnetic-force microscopy images reveal a strong signal for the Cvariants. Since the magnetization of the MFM tip is perpendicular to thesurface, a strong signal indicates an out-of-plane magnetization, whichis parallel to the crystallographic c direction. The c direction is thedirection of easy magnetization in orthorhombic Ni—Mn—Ga martensite[reference 3]. Thus, the conclusion that the C valiant has out-of-planemagnetization is consistent with the magnetic properties of thematerial. The magnetic contrast in A and B variants reveals magneticdomains which are elongated along the crystallographic c direction. Thewidth of the magnetic domains is the same for A and B variants. Thisindicates that magnetic domains have a rod-like shape. The domains matchacross twin boundaries (FIG. 8 b) to ensure magnetic compatibility byavoiding magnetic charges on the twin boundaries [reference 22]. Similarmagnetic domain structures were reported for tetragonal Ni—Mn—Gamartensite [references 23-25].

The results of the indentation experiments show that after unloadingonly about 30 percent of the total deformation remains. Permanent(plastic) deformation may be due to twinning or dislocation activity.Twinning causes reorientation of the lattice. In the A and B variants,twinning would produce C variants. C variants can be identified by MFMimaging via their strong signal stemming from out-of-plane magnetization(FIG. 8 a). For most of the indents, the MFM signal did not displayedvariation compared with the surrounding area. This indicates that theresidual deformation was carried by lattice dislocations. This is aninteresting result since in uniaxial compression tests, twinning occursin this material at only a few MPa stress [references 4, 5] whereasyielding by dislocation motion requires a much higher stress level[reference 26]. Furthermore, under uniform loading, Heusler alloys tendto fail by brittle fracture [reference 26].

In nanoindentation experiments, the applied average stress under the tiparea is largest at the beginning of the indentation since the area isvery small. At a 10 nm indentation depth, the average normal stress atthe surface is about 10 GPa (FIG. 8 b). This is an enormous stresslevel. Thus, the stress near the surface might be sufficient to grow atwin. However, the stress field extends into the material only by aboutthree times the indentation depth and decays very rapidly beyond that[reference 27]. Therefore, if a twin forms, it can grow only to a lengthof about three times the indentation depth. At this short distance, theinteraction of the twinning disconnections with the surface is stillstrong. Upon unloading, the image force causes detwinning (i.e.pseudo-elastic twinning). This assumption follows Garber's conclusion(see discussion in reference [18]) that pseudo-elastic twinning requiresstress concentrations while homogeneous loading leads to permanenttwinning. Following this line of thought, the stress distribution causedby the accidental impact by the AFM covered a large enough area to causepermanent (plastic) twinning. One might conclude that with increasingindenter tip size and tip radius, there should be a transition frompseudo-elastic twinning to plastic twinning. Tall et al. [reference 28]indeed reported indentation-induced permanent twinning in LaAlO₃ using amicroindenter. This indicates a size-effect on the deformationmechanisms operating under localized loading, which might be investigateusing tips with different tip radii [reference 29].

About 70 percent of the deformation is recovered. It is unlikely,however, that 70 percent of the deformation is purely elastic. Therecovery of deformation may stem from pseudo-elastic twinning or fromthe recovery of dislocations due to the image force. Dislocationscontribute at least partially to deformation. This follows from highlylocalized deformation in the pileups next to the indents, which by farexceeds the twinning shear. Thus, dislocation motion is required to formpileups. However, the intensive and irreversible mutual interaction ofdislocations prevents full recovery of the deformation after unloading[reference 30]. Thus, the large recoverable deformation is an indicationthat pseudo-elastic twinning took place.

6. Conclusions

-   -   The crystallography and spatial configuration of twinning        systems can be characterized by using AFM provided a limited        knowledge about the orientation state of one martensite variant.    -   Different twinning modes are active in orthorhombic Ni—Mn—Ga        martensite. The ab- and bc-twinning modes were identified.    -   Magnetic domains in orthorhombic Ni—Mn—Ga martensite have the        shape of rods elongated along the crystallographic c direction.    -   C variants are magnetized perpendicular to the surface and cause        a strong stray-field leading to strong interaction with an MFM        tip.    -   After nanoindentation about 70 percent of the deformation        recovers. The residual deformation is carried by dislocation        activity.    -   It is likely that pseudo-elastic twinning takes place during        nanoindentation.    -   The results indicate a size-effect on twinning deformation under        localized loading.

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Some embodiments may be described as follows: The preferred systemcontrols and detects the crystallographic state of a non-cubic crystal.Said crystallographic state may be connected/related to a magnetic stateand/or to a shape state. Said crystallographic state may be controlledby a magnetic field and/or by a mechanical force field and/or an A/Celectric field. The crystal may be a magnetic shape-memory alloy. Thecrystal may be magnetically and structurally anisotropic, and themagnetic and crystallographic anisotropy coupled. Said magnetic statemay be controlled by a mechanical action. The magnetic state and/or thecrystallographic state may be detected with the magneto-optical Kerreffect. The magnetic state and/or the crystallographic state may bedetected with spin-polarized electrons. The magnetic state and/or thecrystallographic state may be detected with optical interferometry, orwith other or a combination of methods/effects. A system may be providedwherein two different states/functionalities selected from said magneticstates, crystallographic states, and shape states are simultaneouslyused as read and write operators. Two different states/functionalitiesmay be selected from said magnetic state, crystallographic state, andshape state, said two different states/functionalities that are selectedmay be simultaneously used as sense and indicate operators. Twodifferent states/functionalities selected from said magnetic state,crystallographic state, and shape state may be simultaneously used assense and control operators. Said magnetic states and/orcrystallographic states and/or shape states may be embodied in crystalattached to or grown on a substrate. Said magnetic state and/orcrystallographic state and/or shape states may be embodied in individualbits/elements embedded in a crystal.

In some embodiments, a system for writing and reading data comprisescontrol and detection of crystallographic states of a magneticshape-memory alloy or other magnetoplastic and/or magnetoelasticmaterial. Said crystallographic state may be connected/related to amagnetic state and/or to a shape state. Said crystallographic state maybe controlled by a magnetic field and/or by a mechanical force fieldand/or an A/C electric field. The crystal may be magnetically andstructurally anisotropic and the magnetic and crystallographicanisotropy are coupled. Said magnetic state may be controlled by amechanical action. The magnetic state and/or the crystallographic statemay be detected with the magneto-optical Kerr effect. The magnetic stateand/or the crystallographic state may be detected with thespin-polarized electrons. The magnetic state and/or the crystallographicstate may be detected with optical interferometry. The magnetic stateand/or the crystallographic state may be detected with other methodsthan mentioned above. The magnetic state and/or the crystallographicstate may be detected with a plurality/combination of effects selectedfrom the group consisting of: magneto-optical Kerr effect,spin-polarized electrons, and interferometry. The magnetic state and/orthe crystallographic state is detected with a plurality/combination ofeffects selected from the group consisting of: magneto-optical Kerreffect, spin-polarized electrons, interferometry, and other effects notincluding diffraction-based methods. The system may be further adaptedfor sense&indicate functionality, and/or further adapted forsense&control functionality. Embodiments of the invention may comprisemethods of writing and reading data/information comprising use of any ofthe systems described herein; and/or methods of sensing and indicatingdata/information comprising use of any of the systems described herein.In preferred embodiments of the apparatus, systems, and methods, saidcrystal comprises at least six crystallographic states. Preferably, saidcrystal further comprises multiple magnetic states.

Although this invention has been described above with reference toparticular means, materials, and embodiments, it is to be understoodthat the invention is not limited to these disclosed particulars, butextends instead to all equivalents within the broad scope of thefollowing claims.

1. A logic or memory device comprising a magnetoplastic and/ormagnetoelastic material having at least six states.
 2. A device as inclaim 1 which is a combined logic and memory device.
 3. A device as inclaim 1 wherein the memory is nonvolatile.
 4. A device as in claim 2wherein the memory is nonvolatile.
 5. A device as in claim 1 whereinsaid states are shape states.
 6. A device as in claim 1 wherein saidstates are magnetic states.
 7. A device as in claim 1, wherein saidmaterial is magnetic shape memory alloy.
 8. A device as in claim 1,wherein said material is selected from the group consisting of: Heuslershape-memory alloy, non-Heusler ferromagnetic shape-memory alloy,antiferromagnetic shape-memory alloy, non-shape-memory magnetoplasticalloy, Ni₂MnGa, Dy, α-Fe, Co—Ni, τ-MnAl—C, L1₀FePd, L1₀CoPt, L1₀FePt,Ni₂MnGa, Co₂NiGa, Ni₂MnAl, Ni₂FeGa, Fe₃Pd, Fe—Pd—PT, γ-Mn—Fe—Cu,Ni₅₁Mn₂₈Ga₂₁, Ni—Mn—Ga alloys, Ni—Mn—Ga ferromagnetic martensite, andcombinations thereof.
 9. A method of writing data on an element havingmore than two states, wherein the element comprises a magnetoplasticand/or magnetoelastic material having at least six states.
 10. A methodas in claim 9, wherein the writing step comprises application of amagnetic field change.
 11. A method as in claim 9, wherein the writingstep comprises application of a mechanical force.
 12. A method as inclaim 9, wherein the states are shape states.
 13. A method as in claim 9wherein said states are magnetic states.
 14. A method as in claim 9,wherein said material is a magnetic shape memory alloy.
 15. A method ofreading data from an element having more than two states, wherein theelement comprises a magnetoplastic and/or magnetoelastic material havingat least six states.
 16. A method as in claim 15, wherein the writingstep comprises application of a magnetic field change.
 17. A method asin claim 15, wherein the writing step comprises application of amechanical force.
 18. A method as in claim 15, wherein the states areshape states.
 19. A method as in claim 15 wherein said states aremagnetic states.
 20. A method as in claim 15, wherein said material is amagnetic shape memory alloy.
 21. A method as in claim 15, wherein thereading step is conducted with atomic-force microscopy.
 22. A method asin claim 15, wherein the reading step is conducted with magnetic-forcemicroscopy.
 23. A method as in claim 15, wherein the reading step isconducted with spin-polarized electrons.
 24. A method as in claim 15,wherein the reading step is conducted with magneto-optical Kerr-effect.25. A method as in claim 15, wherein the reading step is conducted withoptical inferometry.
 26. A method of writing data on, and reading datafrom, an element having more than two states, wherein the elementcomprises a magnetoplastic and/or magnetoelastic material having atleast six crystallographic states.
 27. A method of claim 26 wherein thewriting step is selected from the group consisting of application of amagnetic field and application of mechanical force.
 28. A method as inclaim 26 wherein the reading step is selected from the group consistingof: atomic-force microscopy, magnetic-force microscopy, spin-polarizedelectrons, magneto-optical Kerr effect, and optical interferometry. 29.A method of combining memory and logic functionalities in a singleelement having more than two states, wherein the element comprises atleast one magnetoplastic and/or magnetoelastic material having at leastsix states.
 30. A method as in claim 29, which combines read and writecapabilities.
 31. A method as in claim 29, which combines sense andindicate capabilities.
 32. A method as in claim 29, which combines readand control capabilities.
 33. A method as in claim 29, wherein said atleast six states are shape states.
 34. A method as in claim 29, whereinsaid at least six states are magnetic states.
 35. A device as in claim 1that is a data reading and writing device.
 36. A device as in claim 1that is a displacement and sensing device.
 37. A device as in claim 1that is a sensing and control device.
 38. A device as in claim 1 that isa sensing and indicating device.